Process for treating austenite stainless steels



Oct. 13, 1964 A. LULA ETAL PROCESS FOR TREATING AUSTENITE STAINLESS STEELS Filed Oct. 3, 1962 3 Sheets-Sheet 1 Austenite Ausfeniie Ferrite Ausieni'e Mnrtensite Feniie Oct. 13, 1964 R. A. LULA ETAL 3,152,934

PROCESS FOR TREATING AUSTENITE STAINLESS STEELS 3 Sheets-Sheet 2 Filed Oct. 3, 1962 Oct. 13, 1964 R. A. LULA ETAL 3,152,934

PROCESS FOR TREATING AUSTENITE STAINLESS STEELS 5 Sheets-Sheet 3 Filed 00 5, .1962

ig. IO

. itation hardening mechanism imposed thereon.

treatments applied to these steels have resulted in a prod- United States Patent Oflice;

3,l52,934 Patented Qct. 13, 1964 3,152,934 PRGQESS FQR TREATING AUSTENHTE STAINLEdS STEELS Remus A. Lula and doseph A. Ferree, In, i atrona Heights, and Thomas H. Mcilunn, New Kensington,

Pm, assignors to Allegheny Ludlum Steel Corporation,

Breckenridge, Pa., a corporation of Peimsylvania Filed Get. 3, 1962, Ser. No. 228,148 8 Claims. (Cl. 148-136) This invention relates to austentic stainless steels of the transformation hardening and of the transformation and precipitation hardening types, and in particular to a. heat treatment which is suitable for obtaining an optimum combination of mechanical and chemical properties in such steels which are substantially free of delta ferrite in the annealed condition at room temperature.

Heretofore numerous attempts have been made to provide a relatively new family of stainless steels which occupy a position intermediate the conventional AISI Type 300 Series stainless steels and the conventional AlSl Type 400 Series stainless steels. In particular, modern technology has witnessed the need for a stainless steel having characteristics similar to the conventional 300 Series stainless steels as respects the aspects of fabrication and resistance to corrosion and oxidation, yet be amenable to a heat treatment for the development of optimum mechanical properties similar to the conventional 400 Series stainless steels. Because of the particular metallurgical phenomenum involved in the various stainless steels, resort has been had to the development of-transformation hardening stainless steels which mayor may not have a precip- The heat not which has only partially fulfilled the current needs, for either the steel did not develop its full mechanical property potential, or the corrosion resistance became impaired. In order to strike an acceptable balance between these properties, it was found necessary to alter the chemical composition, so that predominantly these steels have exhibited one thing in common, which has become an exceedingly important factor in developing the proper heat treatment for these steels. Through various combinations of balancing of the alloying elements, the resultant steel whether purely transformation hardening or transformation hardening with precipitationhardening mechanism superimposed thereon, exhibited up to about 30% 2 composition, delta ferrite was effectively eliminated from the structure of these steels; however, the heat treatments used for the prior art steels when applied to the steels containing at least about 95% austenite have resulted in a gross failure of the objective in steels. The resulting steel not only displayed far inferior mechanical properties, but also was subjected to a high degree of sensitivity to corrosion.

An object of this invention is to provide a heat treatment suitable for use on transformation hardening type austentic stainless steels substantially free of delta ferrite in the annealed condition at room temperature.

Another object of this invention is to provide a heat treatment for use on transformation hardening austenitic stainless steels with or without a precipitation hardening mechanism superimposed thereon, such austenitic stainless steels being substantially free of delta ferrite, and which will develop optimum mechanical properties without adversely affecting the corrosion and oxidation resistance thereof.

A more specific object of this invention is to provide a suitable heat treatment to be applied to austenitic stainless steels exhibiting, as the major strengthening mechanism, a

transformation hardening phenomenon, and which do not contain more than 5% delta ferrite Within the microstructure in the annealed condition, said steel permitting the use of high carbon contents therein to thereby develop the optimum mechanical properties without adversely affecting the chemical or physical properties of the steels.

Other objects of this invention will become apparent when taken in conjunction with the following descriptions and the drawings inwhich:

FEGURE 1 is a modified Schaei'ller diagram illustrating the balanced relationship of the alloying elements and the effect thereof on the rnicrostructure of the steels to which the heat treatment of the present invention applies;

FIG. 2 is a photomicrograph of a prior art steel illustrating the effect on the microstructure of said steel of the step of high temperature annealing at 1950" F. according to the prior art heat treatment;

FIG. 3 is a photomicrograph of the steel of FIG. 2 after the step of trigger annealing at 1710" F. accordirig to the prior art heat treatment process;

FlG. 4 is a photomicrograph-of the steel af FIG. 3 after having been subjected to a sub-Zero cooling treatment at 100 F. and a tempering treatment at 850 F. in

' accordance with the prior art practice;

ofydelta ferrite within the microstructure of the steel.

- at supersonic speeds.

Since a compromise between mechanical strength and the chemical properties of the steel was necessary to provide a commercial product, metallurgists and design engineers have expressed the desire for obtaining a similar type of steel having substantially higher mechanical properties.

This is especially desirable where the material'is to be used in.-applications Where a.high strength-to-weight ratio is desired. Since substantially all theprior art steels contain delta ferrite, it was desirable to design a steel substantially free of delta ferrite which would exhibit an out- I standing combination of mechanical properties without adversely affecting the chemical and physical properties of theysteel. 'Through'a modification of the chemical FIG. 5 is a photomicrograph of a steel having a composition within the scope ofthe chemical analysis to which the heat treatment of this invention applies and which has been subjected to a step in the prior art heat treatment, the sameas FIG. 2;

FIG. 6 is a photomicrograph of the steel of FIG. 5 after being subjected to the same treatmentas the steel of PEG. 3;

FIG. 7 is a photomicrograph of the steel of FIG. 6 after having been subjected to the same heat treatment as the steel of FIG. 4;

FIG. 8 is aphotomicrograph of the steel of FIG. 5 after heat treatment at 1950 F- and thereafter cold rolled; f q 1 i FIG. 9 is a photomicrograph of the steel of. FIG. 8 after the steel has been critically'annealed at a temperature of 1850 F.; w 5

PEG. 10 is a photomicrograph of the steel of FIG. 9 after having been subjected to a trigger anneal at a temperature of 1710 F.; and, v a i 7 FIG. 11 is a photomicrograph of the steelof'FIG. '10 after sub-zero cooling and tempering. Each of the photomicrographs of FIGS. 2' through 11 inclusive was taken at a magnification ofSOOX.

In its broader aspects, this invention relates to a process 3 for treating austenitic stainless steels containing at least 95% austenite within the microstructure in the annealed condition at room temperature, said stainless steels being selected from the group consisting of the austenitic stainless steels of the transformation hardening type, either with or without precipitation hardening mechanism superimposed thereon. These steels have a composition encompassing a series of alloying elements within the range of up to about 0.40% carbon, up to about 8.0% manganese, up to 2.0% silicon, from about 8.0% to about 20.0% chromium, up to about 13.0% nickel, up to about 4.0% molybdenum, up to about 0.60% nitrogen, up to about 4.0% of metal selected from the group consisting of aluminum, copper and vanadium, and the balance iron wtih the incidental impurities usually found in the normal steel-making practice.

Reference is directed to Table I which illustrates the general range and the preferred range of alloying elements which are especially suitable for the heat treat- 1 Metal selected from the group A1, V and Cu.

The steel so selected as to have a composition within the limits set forth hereinbefore in Table I may be manufactured in any of the well-known steel-making processes, the steps of which are well known in the art and need not be set forth in detail herein. It is sufficient to relate that a melt having an analysis within the general range set forth in Table I and which is balanced to have a chemical composition to produce an austenitic stainless steel exhibiting at least 95% austenite in the annealed condition at room temperature is teemed into ingots which, when solidified, may be hot worked in any desirable manner. Preferably the ingot is hot rolled into the form of a strip which is coiled in the usual manner into a product commonly called a hot rolled band. The hot rolled band is thereafter pickled and annealed. The annealing treatment is conducted usually on a continuous basis and hot rolled band, which is to be subsequently cold rolled, is annealed at a temperature in excess of 1900 F. which is effective for producing the steel in its most stable form, that is, substantially all of the carbides are taken into solution, thus stabilizing the austenite phase of the steel by reducing the M temperature to substantially below room temperature. This is especially useful from the standpoint of cold rolling since the annealed hot rolled band can be reduced in cross sectional area to a greater degree before reannealing becomes necessary, because these austenitic stainless steels undergo a considerable amount of work hardening during cold working operations. Subsequent to the cold rolling operation, that is, after the hot rolled band has been reduced in cross sectional area to the finish gauge, it is again reannealed at a temperature in the range between 1800 F. and 1875 F. prior to shipment to the fabricator or ultimate user. Thetemperature to which the steel is heated subsequent to the cold forming of the steel to the finish gauge mustnot in any event exceed the limit of the carbon solubility in the steel, which, in the steels to which this heat. treatment applies, is usually about 1875 F. maximum.

- Reference is now directed to FIG. 1 which is a modified Schaefiler diagram illustrating a pseudo-phase diagram relationship between the austenitizi'ng elements and the ferritizing elements. It is to be noted that FIG. 1

more closely approaches equilibrium conditions than the original Schaeffler diagram which was based upon cast welded material. In this respect, it will be noted that the phase boundary lines have been transposed downwardly and to the left from the position that they originally were illustrated in the original Schaeffler diagram. It must also be pointed out that these phase relationships were based on the following data in which a constant silicon and manganese content of 0.40% and 1.0%, respectively, were used along with the actual percent nickel and the nickel equivalents and the actual percent chromium and the chromium equivalents. Specifically, the relationship between nickel and the nickel equivalents is expressed as follows:

Ni equivalents=percent Ni+30 (percent C -l-percent NH- /z (percent Mn)+0.25 (percent Cu) and the relationship between chromium and chromium equivalents Cr equivalents=percent Cr+ percent Mo +1.5 (percent Si) +2 (percent Al) +2 (percent V) It will be appreciated that the nickel equivalents eX- press the relationship of the relative power of the austenitizing elements and their effect upon the microstructure of the steel, whereas the chromium equivalents express the relationship of the power of these elements in their ferritizing effect upon the steel.

By inspection of FIG. 1, single phase and twoand three-phase areas are apparent. While a steel having a composition within the limits set forth hereinbefore in Table I and which has a chromium equivalent and a nickel equivalent as computed in accordance with the equations given hereinbefore may fall within the limits of the partial phase diagram of FIG. 1, it is only the steels which have a balanced composition of nickel equivalents and chromium equivalents which fall generally within the portion designated ABCDEFGA and preferably within the area aBCdEFGa or the portion ABCDEGA which will respond to the heat treatment to be set forth more fully hereinafter to provide the steel with an optimum combination of the attainable mechanical properties. It will be noted that the area bounded by ABCDEFGA in FIG. 1 will normally have a-microstructure at room temperature in the annealed condition which will exhibit at least austenite, the balance thereof may be either ferrite or martensite. Preferably, the steels to which the process of this invention applies will have a composition which is preferably confined to the all austenite field aBCdEFGa as more particularly set'forth in FIG. 1. Thus, it will be seen that while it ispossible to have a composition substantially completely austenitic at room temperature in the annealed condition, there may also be present small amounts ofdelta ferrite, the total delta ferrite content never exceeding 5% of the microstructure, or in the alternative, the steel may also contain martensite, the martensitic constituent preferably being less than about 10%. It is these steels and only these steels which have a composition falling within the area ABCDEFGA and preferably within the area aBCdEFGa and which are subject to hardening by a transformation hardening mechanism, and either with or without an age hardening mechanism superimposed thereon, which will respond to the process of this invention to provide an optimum combination of mechanical properties in the steel.

As was stated hereinbefore, the prior art steels struck an acceptable balance between the mechanical and chemical properties through a modification of the chemical analysis by balancing the'alloying components to such a degree as to produce a steel having a microstructure which exhibited about 10% to 30% delta ferrite therein. These prior art steels were normally annealed at temperatures in excess of 1900" F. and preferably at a temperature within the range of between about 1950 F. and

2050" F.; however, at these temperatures, the steels have the result that if the ultimate grain size is not controlled there will be an adverse effect on the mechanical properties of the steel. Consequently, it was found that with a minimum of about delta ferrite present within the microstructure of the steel and distributed randomly throughout, grain coarsening was efiectively inhibited to sucha degree that'the high annealing temperatures could be used without excessive grain growth occurring in said steels. Thus, with between 10% and 30% delta ferrite present within these prior art steels, they can be annealed at temperatures in excess of 1)00 F. without the steel exhibiting a grain size larger than about ASTM No. 7.

Reference is directed to the drawings and FIGS. 2 and 5 in particular. FIG. 2 illustrates the microstructure of Heat No. 79650, a steel which has a chemical composition as set forth in Table II hereinafter and which is within the General Range of Table I, but is so balanced as to have a microstructure containing more than 10% delta ferrite therein. Contrasted to this, the steel of FIG. 5, which is Heat No. 79876 having a composition as set forth in Table II hereinafter, while being within the General Range of Table I, has a balanced composition so that it is substantially free of delta ferrite. FIG. 2 illustrates the effect of the high annealing temperature, that is, in excess of 1900 F., in this particular instance 1950" E, on a steel containing more than 10% delta ferrite in the microstructure, whereas FIG. 5 illustrates the effect of the same annealing treatment on a steel substantially free of delta ferrite. It is apparent that the delta ferrite 10 of the steel of FIG. 2 is randomly distributed throughout the austenitic matrix 12. It is to be noted in particular that grain growth has been effectively inhibited by the presence and random distribution of the delta ferrite 10, the actual grain size rating being ASTM No. 7 or smaller. Contrasted to this, it will be seen in FIG. 5 that excessive grain growth of the austenite 14 has taken place in this steel. It will be appreciated in FlG. 5 that the steel after annealing at a temperature of 1950 F. was cold rolled 20% in order to partially transform some of the austenite to martensite. It was found necessary to do this from a purely metallographic point of view so that some contrast would be afforded in the photomicrograph to identify the grain boundaries, it being noted that all attempts to provide an etching reagent to outline the grains failed.

It is possible to use other means to illustrate austenitic grain size such as were used in producing FIG. 6 hereof. The means include a low temperature heat treatment to purposely precipitate grain boundary carbides which outline the austenite grains as illustrated in FIG. 6. As illustrated in FIG. 5, which was taken at a magnification of 500x, the austenite exhibits a very coarse grain size. When said grain size was rated according to ASTM standards, the steel exhibited a grain size of ASTM No. 45 or larger. Thus it is easily seen that the delta ferrite has performed its first important expected since the temperature of 1950 F. is above the limit of the carbide solubility and therefore substantially all carbides are taken within solution with the austenitic matrix. Annealing at temperatures in excess of 1900 F.

V is effective for taking most, if not all/of the alloyin 5 elements, especially the carbides, within solution, either 1 within the austenite or within the delta ferrite. As a result, the steel is in its most stable condition for that particular chemical composition as respects theiniiuence or heat treatment on the steel. In other words, the M temperature of the steel is depressed to a temperature substantially below room temperature. 'It is for this reason that it is desirable to use high annealing temperatures, that is, in excess of 1900 F., because since the steel is l 1 in its most" stable form, it will permit amyriad offabri- 6 cation operations to be performed on the steel without the probable transformation of the steel during these operations, and in some cases without the necessity of reannealing the steel betweenfabrication operations.

The prior art steels after annealing at temperatures in excess of l900 F. do not possess the optimum mechanical properties; therefore, it becomes essentialto heat treat the steel to obtain the desired combination of both mechanical and chemical properties. Resort was had to reannealing these steels at a temperature in the range between about 1650 F. and about 1800 F. which was effective for purposely precipitating a portion of the carbides which had been taken into solution at the high annealing heat treatment temperatures. It was necessary to precipitate a portion of the carbides to unbalance the austenite phase to thus raise the M temperature to about room temperature or some conveniently fixed lower temperature so that the subsequent treatments could produce the desired transformation phenomenon, thereby obtaining the desired mechanical properties Within the steel. It was during this unbalancing of the austenite that the second vital function of delta ferrite came into play.

When the steels containing at least 10% delta ferrite were reannealed at a temperature in the range between 1650" F. and about 1800 B, it Was found that the carbides which precipitated during this heat treatment appeared at the austenite-delta ferrite grain boundary interface. This resulted in the austenite-austenite interface being relatively free of carbide precipitates.

It will be appreciated that since the delta ferrite is discontinuous, that is, it is randomly distributed throughout the austenite matrix and it is along the austenite-delta ferrite interfaces that the carbides have precipitated, there is no continuous area which is depleted of chromium; hence the corrosion resistance is not measurably impaired. As stated, while the precipitation of the carbide did not have too pronounced aneifeet on the corrosion resistance properties, the austenite was depleted of a portion of the carbon content with the result that the M tempera-. ture was raised to about room temperature or above, thus permitting the steel to be transformed to martensiite during the subsequent treatments.

As demonstrated previously, an additional characteristic of vthese steels containing at least 10% delta ferrite is that the grain size is effectively inhibited from excessive coarsening during the high temperature anneal. Since the carbide precipitation takes place at preferred sites during the l650 F.1800 F. trigger anneal, in this case the ferrite-austenite interfaces, it is believed that car bon migration or diffusion from the central areas of each grain to the boundaries takes place and is an essential part of the transformation phenomenon since this is the mechanism by which the M temperature is raised. With fine grain sizes, it appears that the grains become, more or less uniformly depleted of carbon within the time allowed at temperature. Consequently, there is no significant carbon gradient across the individual grains, the M temperature is uniformly raised, and uniform, complete transformation of the austenite to martensite is obtained in subsequent treatments. The ferrite, of course,

erties are concerned, and the average mechanical properties of the steel compared to the maximum possible with a substantially completely martensitic structure, are reduced in proportion to the amount of ferrite present.

On the other hand, the steel to which the process of thisinvention applies does not react in the same manner as a steel containing delta ferrite. The substantially completely austenitic' stainless steel which has been annealed at a temperature in excess of l900 F. and whichwas not cold worked subsequent thereto, was subjected toa re annealing 2 between 1650 F. and 1800 F. to purposely precipitate a portion of the carbides from solution'in order, to uneat treatment'at a temperature in the range cipitated, formed at the austenite ustenite grain boundary interface, thus forming a substantially complete envelope of grain boundary carbide precipitate. The effect of the precipitation of the grain boundary carbides will be discussed more fully hereinafter.

Reference is directed to FIGS. 3 and 6 of the drawings which illustrate the effect of reannealing the steels of FIGS. 2 and 5, respectively, at a temperature of 1710 F. In the steel of FIG. 3 it is apparent that the chromium carbides 20 are found at the interface 22 between the delta ferrite 1t and the austenite 12. The austenite-toaustenite interface 24 is relatively free of the carbides 2% in this steel. On the other hand, it is readily apparent that there is no delta ferrite present within the microstructure of the steel of FIG. 6. Since the annealing temperature is below the limit of carbon solubility, a portion of the carbides must precipitate, and as illustrated in FIG. 6, they have precipitated at the austenitesaustenite interface 24. This substantially continuous envelope of grain boundary carbide precipitates combined with a large grain size is highly deleterious to the mechanical properties of the steel.

It is believed that the development of these grain boundary carbide precipitates also produces another significant effect which highly influences the mechanical properties. Since the carbides form at the grain boundaries, thus depleting the area adjacent the grain boundary interface of a portion of its carbon content, there is, in effect, a micro-gradient in the chemical composition across the grain. This micro-gradient effect becomes more pronounced with increasing grain size.

It is believed that because of the large grain size, the carbon migration or dilfusion is not sufficient to substantially uniformly deplete the grain of carbon within the time period that these steels are maintained at the elevated temperature. Since a uniform depletion of the carbon content across the grain is an essential part of the transformation of the phenomenon because it is the mechanism by which the M temperature is raised, [the large grain size makes it virtually impossible to obtain a substantially uniform depletion of the carbon content. Thus, the M temperature is not uniformly raised with the result that there is not a uniform complete transformation of the austenite to martensite obtained at the subsequent heat treatments. The end result is that inferior mechanical properties are developed in the steel. The adverse effect on the mechanical properties is realized in the following manner: the subsequent hardening treatment to which the steel is subjected usually consists of a sub-zero cooling treatment followed by a tempering treatment. The effect of the sub-zero cooling treatment is to transform the austenite to martensite and of course the function of the tempering treatment is to temper the martensite and effect some degree of aging. However, when the microstructure of the steel exhibits a continuous envelope of grain boundary carbides and a large grain size, it has been found that the interior of the grain is quite stable so that when the steel is sub-zero cooled to transform the austenite to martensite, the martensite constituent is because of the distribution of the austenite, which is caused by the concentration gradients within the grain produced by the grain boundary carbide precipitation, poor mechanical properties result. It will be appreciated that in these transformation hardening type austenitic stainless steels, the microstructure in the hardened condi-.

tion exhibits between about and about 25% retained austenite. Thisis true regardless of the heat treatment applied to these steels. However, it is believed that'the distributionof the retained austenite may be a significant factor in. some steels exhibiting poor mechanical properties. Thus, when the retained austenite is concentrated in large areas near the center of the grain, for example, as will be shown hereinafter with respect to FIG. 7, the steels exhibit poor mechanical properties, whereas if the austenite is substantially uniformly distributed throughout the grain, a significant improvement in mechanical properties is obtained. As will be more clearly demonstrated hereinafter, the steels, which contain large amounts of retained austenite which is substantially concentrated in the central portion of the grain, exhibit far inferior mechanical strength properties which, by comparison, are even inferior to those mechanical properties exhibited by a steel containing from 10% to 30% delta ferrite.

Reference is directed to FIGS. 4 and 7 which illustrate the microstructure of the steels of FIGS. 3 and 6, respectively, after sub-zero cooling to a temperature of l00 F. followed by a tempering at a temperature of 850 F. It is readily evident from FIG. 4 that the microstructure consists of the delta ferrite 10 with a substantially completely transformed matrix of martensite 30. On the other hand, FIG. 7 clearly illustrates that the martensite 34) has formed adjacent the grain boundary carbides 22. However, the martensite 30 does not completely traverse the grain, the retained austenite 32 being readily apparent in the central portion of the grain. The effect of the retained austenite on the mechanical properties will be more fully explained hereinafter with respect to Table III.

As is readily apparent from the foregoing, steels containing at least 95% austenite are adversely affected by the use of the prior art heat treatments which are employcd with the steels containing more than 10% delta ferrite. The heat treatment, to be set forth hereinafter, must be applied to steels having a closely controlled chemical composition which exhibit at least 95% anstenite in the microstructure and fall within the area ABCDEFGA and preferably within the area aBCdEFGa of FIG. 1 in order to obtain the optimum mechanical properties. The following heat treatment is especially suitable for the development of optimum mechanical properties in these steels without causing excessive grain growth and without adversely affecting the corrosion or oxidation resistance of the alloy.

In particular, the steel in the form of a hot rolled band may be pickled to remove the mill scale or other objectionable impurities from the surface thereof which may have been picked up during hot working of the steel from ingot into the hot rolled band. While pickling, when necessary, will usually suffice, in some instances it may be necessary to wheelabrate or sand blast the surface to remove tightly adherent defects in the surface quality of the steel. The pickled or otherwise cleaned hot rolled band is next annealed at a temperature dependent upon the subsequent operations to be performed on the steels. Where the steel is to be subsequently cold worked by at least one operation, a temperature in excess of about 1900 F. is used. it is preferred to heat treat the steel at this temperature in order to take all of the carbides within solution, thus stabilizing the austenite phase to its greatest degree. This has the effect of permitting a greater reduction of the cross sectional area of the steel during the subsequent working operations before reannealing becomes necessary.

I As was stated hereinbefore, when these steels are heated to temperatures in excess of 1900 F., excessive grain growth appears to take place. However, at this stage of the processing it is immaterial, unless the grains become so excessively large as to make the steel completely unsuitable. However, before the'grain growth attains such a tremendous size, long-time periods at these elevated temperatures are necessary, and as such they are not economical in commercial production. In commercial operations it is preferred, to anneal the hot rolled band on a continuous basis whereby the hot rolled band is continuously uncoiled and passed through the furnace as is well known in the art. The steel is preferably held at this temperature for a time period ranging between about and 30 minutes and usually a time period only suiticiently great to take substantially all of the carbides within solution. The annealed steel is preferably cooled at a rate sufficiently fast to prevent any carbide pre ipitation. The effect of the annealing heat treatment, in addition to placing substantially all of the carbides Within solution, thus stabilizing the steel, also has the effect of providing the steel with good ductility and low hardness so that a large amount of cold worl: can be applied to the steel without the necessity of reannealing.

The steel in the soft, ductile, annealed condition is next cold Worked in one or more cold worl-jng operations, preferably by cold rolling, to any desired gauge. It will be realized that these steels undergo considerable work hardening during cold rolling and use result, they can only be reduced in cross sectional area to a maximum amount varying between 40% and 70% of the original cross sectional area of the hot rolled band depending upon the type of equipment used. For example, in a 4- ligh reversing cold rolling mill, the steel can be reduced in cross sectional area only about 40%, whereas, it the annealed andpic (led hot rolled band is cold rolled in a Sendzimir mill, the cross sectional area may be reduced as much as 70% before reannealing becomes necessary. When the steel of intermediate gauge has been reduced in cross sectional area to a degree that reannealing becomes necessary, and after reannealing the steel will be subject to further cold working, the cold worked steel of intermediate gauge may be reannealed by again heating to a temperature in the range between about 1960 F. and about 2000 F. at must be pointed out, however, that reannealing at this high temperature is permissible if, and only if, the steel is to be later cold worked to finish gauge and such cold working reduces the cross sectional area at least of the intermediate gauge. Where the cold worked steel has been reduced to finish gauge through the single cold rolling operation, the steel must be reannealed at a temperature in t e range between 1800 F. and l875 F.

The cold worked steel of intermediate gauge which has been reannealed at a temperature in excess of 1900" F. or the limit of carbon solubility, is cold worked to finish gauge. It will be appreciated that the steel in the cold worked condition will not have its best fabrication properties because of the amount of worlc hardening which the steel had undergone during cold rolling to finish gauge. It therefore becomes necessary to reanneal the steel so as to place the steel in its most stable form which is compatible with the subsequent hardening treatmerit to develop the optimum mechanical properties and without adversely atiecting the chemicalproperties of the steel. it is for this reason that the steel of the finish gauge is annealed at a temperature in-the critical range between 1800 F. and 1875" F. This critical annealing temperature range must not be exceeded if the steelis to have the optimum combination of mechanical properties in the final product. An examination of what happens during the final cold rolling andthe critical annealing heat treatment is helpful in understanding the reasons therefor. V

. As was stated, the cold worked steel of intermediate gauge was annealed at a temperature in excess of the carbon solubility limit or 1900 F. in order to place substantially all of the carbides in solution, thus stabilizing the steel as far as possible with respect to the heat treatment. In other words, the annealing heat treatment is effective for depressing the M temperature of the steel to below room temperature. During the subsequent cold working of the steel arter reannealing at above 1909" F., energy is imparted to the austcnite and through the deformation of the steel, the M temperature is raised to above room temperature. As a result, there is a partial transformation of the austenite to martensite. Since the steel is cold worked, at least one of the deformation mechanisms must also be present within the microstructure and in this particular instance the mechanism of slip is readily apparent. In addition, the cold working contributes during the later recrystallization of the steel to provide a relatively small grain size therein. Thus it is seen that the steel, while being in its most stable form after the rcannealing at temperatures in excess of 1900 F., has undergone a partial transformation during cold rolling so that a portion of the austenite grains is transformed to martensite and the balance of the grain is traversed in one or more ways with a number of slip planes, both factors performing Very vital functions as will be more fully described hereinaiter. In particular, it has been found that after steel is cold worked to finish gauge and critically annealed at a temperature in the range between 1800 F. and 1875" F., the steel will exhibit a grain size of about ASTM No. 7 or smaller.

With the cold worked steel of intermediate gauge having been reannealed at a temperature in excess of 1900 F., and subsequently cold worked to eilect a reduction in the cross sectional area of at least 10% to obtain the finish gauge of the steel mill product, the microstructure is again replete with a partial transformation of austenite to martensite and the myriad slip bands traversing the retained austenite. None of the carbides have been precipitated since the steel has been reannealed;

therefore, substantially all carbon is in solution. The partial transformation of the austenite to martensite and the slip bands resulting from cold working are necessary in order to develop the optimum mechanical properties in the steel because we have found that during the critical annealing heat treatment of the cold worked steel of finish gauge, it is within the martensite and along the slip bands that there are preferred sites for the precipi tation or carbide particles. This precipitation is a random precipitation which is in no way concentrated'at the original austeuite-austenite grain boundary interface.

Upon heating of the cold rolled steel of finished gauge from room temperature up to the critical annealing temperature range, a number of processes occur which are,

significant to the process of this invention. Upon the application of heat to the cold rolled steel, the martensite constituent of the microstructure is tempered. At the same time, carbide particles begin to precipitate randomly throughout the martensite. Simultaneously, carbide particles will precipitate along the slip bands and in the retained austenite, said carbide precipitate being randomly distributed with respect to the microstructure. As heating continues to the temperature of about 1500 F., the critical temperature is surpassed, at which time the martensite retransforms to austenite. The carbides continue to precipitate at an increasing rate within the austenite until the temperature of about 1500 F. is attained. Moreover, while the martensite is retransforming to austenite the processes of recovery and recrystallization are also taking place, thus destroying the slip bands which were apparent just after cold rolling. As the temperature is raised to the critical annealing temperature, that is, a temperature in the range between 1800 F. and in no case greater than 1875 F, the car-. bides which have been precipitated in a random distribu tion along the slip planes, martensite and retained austenite are partially redissolved within the austenite. Since the temperature to which the steel of finished gauge is heatedduring the critical anneal is neverin excess of 1875 F., not all of the carbides are taken in solution retain its relatively small grain size in comparison with a steel heated to a temperature in excess of 1900 F. The fact that part of the carbide particles remain out of solution becomes very important as will be more fully explained hereinafter with respect to the trigger anneal of the steel during the hardening treatment.

Reference is directed to PEG. 8 which illustrates the effect of cold rolling the steel subsequent to the intermediate reanneal at a temperature in excess of 1900 F. The microstructure of the steel of FIG. 8, which is Heat No. 35558 having a composition as set forth in Table II hereinafter, clearly shows austenite 40 having a portion of each grain partially transformed to martensite 42. It will be appreciated that the microstructure of the steel as illustrated in FIG. 8 does not visibly exhibit slip bands as would be expected from what was said hereinbefore. However, while the slip bands are present, the metallographic polishing of the steel removes the visible evidence thereof, hence their absence from FIG. .8. Note in particular that there are few carbides apparent. Thus, it is clear that the effect of the intermediate anneal at a temperature in excess of 1900 F. is sufiicient to place all carbides within solution within the austenite. The steel is illustrated in FIG. 8 when reannealed to the critical annealing temperature, that is, to a temperature within the range of 1800 F. to 187S F., exhibits a microstructure as shown in FIG. 9. As illustrated in FIG. 9, the microstructure consists of austenite 50 having randomly distributed carbides 52 dis tributed throughout the grain. Thus, it will be seen that it is of primary importance that the steel which has been cold worked to effect a reduction of area of at least 10% toThe finish gauge, be reannealed below the carbon solubility limit and preferably at a temperature in the range of 1800 F. and 1875 F. in order to prevent the grain boundary carbide precipitation during the hardening heat treatment as illustrated in FIG. 6.

The elfect of annealing the steel at this critical temperature is to provide the steel in its most stable form compatible with optimum response to the treatment to be described hereinafter to develop the optimum mechanical properties without adversely affecting the corrosion or oxidation resistance. It is desired to maintain the steel in its stable form because the steel can be formed into its finished shape without the necessity of reannealing. Another consideration to be noted is the fact that the fundamental hardening pehnomenon exhibited by this steel is basically transformation hardening mechanism involving a transformation of austenite to martensite. As a result, it is desired to have the steel in its most stable form for shipment because, with the M temperature substantially below room temperature there is little, if any, tendency for transformation in transit, whereas, if the steel is not in its most stable form, cooling to below room temperature may be sufiicient for transforming a portion of the austenite to martensite with the result that the steel will \have poor fabrication properties.

After the steel has been fabricated to its finished shape, and in some cases before, where the fabrication to be done on the steel is not too severe, the cold rolled steel of finish gauge which has been given the critical anneal at a temperature in the range between 1800 F. and 1875 'F. is subjected to the hardening heat treatment. This consists of subjecting the steel to a trigger anneal to purposely unbalance the austenite phase through the precipitation of a predetermined amount of carbide particles randomly throughout the microstructure of the steel. It is preferred to heat the steel to a temperature in the range between about 1650 F. and about 1750 F. and preferably at a temperature in the range between 1685 F. and 1735 F. It will be noted that whether or not the steel is subjected to some type of deformation during fabrication, or whether the steel is to be hardened without deformation subsequent to the critical anneal, maximum response of the steel to the hardening treatment requires the steel to be Cir subjected to the trigger anneal in order to develop the optimum combination of mechanical properties within the steel. It is believed that the critical anneal at the temperature within the range of 1800 F. to 1875 F. may create'a substantial equilibrium at this temperature so that the carbon saturated austenite at this temperature will not be sfiicicnt to take all of the carbide particles in solution. Upon heating to the trigger annealing temperature range, this equilibrium is disturbed so that, in effect, the austenite phase becomes supersaturated with carbon at this temperature. Consequently, the reaction will reverse and carbides will be precipitated. It is for this fundamental reason that it is necessary, in fact absolutely essential, that the annealing temperature of the critical anneal does not exceed the carbon solubility limit. Since a portion of the randomly distributed precipitated carbides stil remain within the austenitic phase of the microstructure during the critical anneal, the trigger anneal is effective for precipitating a portion of the carbides at these preferred sites, thus unbalancing the austenite by raising the M temperature to room temperature or above. In thus rejecting the carbon from the austenite phase, it precipitates as ohrornium carbide at the preferred sites of the presently existing carbide precipitates. This, in effect, produces a substantially random distribution of the carbides throughout the austenite grains.

Reference may be had to FIG. 10 which illustrates the eifect of annealing the steel of FIG. 9 at a temperature of 1710 F. It is immediately apparent that there is a great increase in the amount of carbides 52 which have been precipitated in comparison with the steel of FIG. 9. Note in particular that the distribution is random and very little grain boundary carbides appear, it being noted that whatever grain boundary carbides are present do not form a substantially continuous envelope of grain boundary precipitates about the grain.

Following the trigger anneal, the steel is next subjected to a sub-zero cooling treatment, usually at a temperature in the range between F. and 1l0 F. While lower temperatures can be used, these temperatures are usually sufiicient to attain a substantially completely martensitic structure. Cooling to these temperatures is sufficient to transform the austenite to martensite, the hard phase of steel. It should be pointed out in particular that due to the chemical composition of the steels substantially free from delta ferrite, 'a higher carbon martensite may be formed. As is known, the carbon content of the martensite determines the ultimate hardness of the martensite within relative limits. Therefore, since the steels to which the process of this invention applies usually contain higher carbon contents than the steels containing delta ferrite which have compromised the chemical composition in order to obtain an acceptable balance between mechanical properties and chemical properties, this heat treatment is effective for producing the optimum mechanical properties in steels of substantially higher carbon contents and without adversely affecting the chemical properties thereof. The martensite thus formed will exhibit its full mechanical property potential in the finished product. Subsequent to the subzero cooling treatment, it is preferred to temper the steel in order to obtain the optimum balance of mechanical properties, that is, between strength, ductility and hardness. While the steel in the martensitic condition exhibits outstanding strength and hardness, the yield strength is low. Consequently, the eifect of tempering treatment is to substantially increase the yield strength without adversely affecting the tensile strength or hardness to any considerable degree. The tempering treatment is preferably performed by heating the steel to a temperature in the range between 750 F. and 1050 F. for a time period of up to about 16 hours and thereafter air cooling. This treatment may also be effective for precipitation hardening the steel. In some instances, it may be necessary to double temper the steel in order to take '13 full advantage of the precipitation hardening mechanism which is superimposed on the transformation hardening mechanism. Reference is directed to FIG. 11 which 11- lustrates the effect of the sub-zero cooling andtempering treatment on the steel of FIG. 10 and clearly illustrates cept FG-lZ, FG11 and FG-l were hardened by the following treatment, viz., trigger annealing at a temperature of 1710 F. followed a quench to room temperature and thereafter they were sub-Zero cooled to 100 F. for three hours followed by a tempering treatment at 850 F. for 3 hours. Metallographic examination revealed that all steels except Heat Nos. 79148, 35204 and FG-12 were substantially free of delta ferrite.

Referring now to the test data set forth in Table Hi, and in particular to the test results recorded for Heat 7965 0, which has at least 10% delta ferrite in the microstructure at room temperature, it can be seen that the an- TABLE II Chemzcal Composztzon (Percent by Wt.)

4 115 1100. Eggs, 0 Mn s1 01 N1 Mo N Al Cu Fe or. N1,

0.089 0.76 0.24 10.37 4.11 2.78 Balance. 10.40 9.83

Most of the steels set iorth in Table ii are commercial heats, a portion of which were processed, both by prior art methods and the method set forth in this invention.

These steels with their respecnve annealing heat treat ment as set forth hereinafter 1n Table HI were thereafter hardened as set forth and then the steels were glven the standard tens1le test which compromises the measuring of the tensile strength, yield strength and-ductility, the latter as illustrated by the percentage elongation and the standard hardness test.

TABLE III T ensile Proper-mes Annealing Ultimate Yield Elong Heat N0. Temp. Tensile Strength i112 Hardness F.) Strength (p.s.i.) (percent) 11,,

(p.s.i.)

1,800 216,000 184,100 0.0 1, 850 222, 000 187, 000 11. 5 1, 875 222, 000 187, 000 11. 5 1. 000 210, 000 182, 000 10. 5 1, 025 217, 000 179, 000 9. 0 a 1, 050 210, 000 170, 400 70201 1,750 214,000 0.0 1, 800 215, 100 187, 600 11. 5 1, 850 221,000 185,000 0. 0 1, 875 220,000 185, 000 10. 0 1, 025 215, 000 181, 000 10. 5 1, 050 202, 000 161, 300 11. 0 0x035 1,850 239, 750 104,350 12.0 1, 000 240, 250 104, 150 13. 75 1, 050 228, 780 176, 170 12.75 2,000 218, 860 158. 710 10. 75 FG12 1,750 220,000 103,000 7.0 1 1,800 225,000 108,000 0.0 1,875 225,000 190,000 0.5 1,050 222,000 180,000 6.5 FG11- 1,800 241,000 212,000 0.0 1, 050 231, 000 100,000 8. 0 -FG-10 1,300 247,000 222,000 0.0 1, 050 233, 000 183, 000' 0. 0 111-21 1,875 250,800 222,500 8.0 1 1,000 222,800 175, 700 11.5

All of the steels set'forth in Table 111 have been cold rolled sufficiently to cause at leasta 10% reduction in the cross sectional area thereof after which they were annealed at the temperatures indicated in Table 111F051- lowing the annealing heat treatment, all of the steels exnealing temperature to which these steels were subjected prior to the hardening treatment has little eifect upon the mechanical properties of these steels. Heat 79650, after annealing at 1800 F., exhibited a tensile strength of about 194,900 p.s.i. and a yield strength of 164,100 p.s.i. When the annealing temperature was increased to' 1875 F, the same steel exhibited a tensile strength of about 199,300 p.s.i. and a yield strength of 164,200 p.s;i. Annealing at a temperature of 1950 F. resulted in the steel exhibiting a tensile strength of 194,300 p.s.i. and a corresponding yield strength of 162,900 p.s.i. Thus, it is clear that the increase of F. in the annealing temperature produced little effect on the mechanical properties of steels containing delta ferrite. From the foregoing test results, it is clear that while the process of this invention is effective for treating steels containing delta ferrite, it does not produce any outstanding etfects on the mechanical properties of these steels as opposed to the mechanicd properties exhibited by these steels when treated by the prior art methods. Reduced to its common denominator, the mechanical properties exhibited by steels containing delta ferrite are independent of the temperature at which they are annealed following cold working to finish gauge.

On the other hand, comparingthe test results recorded for Heats 35526, 35558, 84138 and'7929l, it is immediately seen that these steels possess a substantially higher strength level through a slight modification of the chemical analysis; however, this strength is realized only where the steels are treated according to the process of this invention. Comparing the, test results recorded in Table III from Heat 35526, it is seen that this steel exhibited a tensile strength of.2-l7,200 psi. and a yield strength of 180,300 p.s.i. when the steel was annealed at a temperature of 1875 F. following cold rolling and after annealing was hardened by the treatment set forth hereinbefore. Compared with this level of strength, this same steel when annealed at a temperature of 1900 F., a mere increase of 25 F;, the tensile strength had dropped almost 10,000

p.s.i. and the. yield strength had dropped 14,000 p.s.i. to the values 207,300 p.s.i. and 164,200 p.s.i., respectively. Thus, the critical annealing temperature, the upper limit of which is about 1875 F, clearly exhibits its influence on the mechanical properties of this steel, substantially freeflof delta ferrite. Substantially similar results occur in Heat 35558. 1 The data recorded in Table III for Heat 3555 8 are even more outstanding to illustrate the effect of the critical annealing temperature on the mechanical propstrength of 176,300 p.s.i. when critically annealed at 1850 F. and thereafter hardened. Increasing the temperatureof the critical anneal to 1875 F. has substantially no effect on the mechanical properties. However, when the critical annealing temperature is increased 25 F. to 1900 F. or 50 F. to 1925 F., the change in the mechanical properties is outstanding. The test results recorded in Table III illustrate that the tensile strength has dropped from 213,000 p.s.i. when the steel was critically annealed at 1875 F. to 200,000 p.s.i. and 180,500 p.s.i. when annealed at 1900 F. and 1925 F., respectively. The decrease in the yield strength is even more severe, the test results showing a decrease from 171,800 p.s.i. for a critical annealing temperature of 1875 F. to 153,400 p.s.i. and 122,500 p.s.i. for annealing temperatures of 1900 F. and 1925 F., respectively. Substantially similar results were obtained in the test conducted on Heats 84138 and 79291. In each case, when the steels were critically annealed at temperatures in excess of 1875 F., the tensile properties were adversely afiected.

As was stated hereinbefore, it is through a critical balancing of the alloying elements in these steels so that the composition of the steel falls within the area ABCDEFGA and preferably within the area aBCrlEFGa of FIG. 1 that it is possible to obtain a transformation hardening type stainless steel which is substantially free of delta ferrite. However, it is only through the heat treatment of the process of the invention which makes it possible for these steels, substantially free of delta ferrite, to exhibit outstanding mechanical properties.

Referring again to the test results recorded in Table III, it can be seen by comparing the data recorded for Heat 79650 with the data for Heats 35526 and 35558 when annealed at temperatures in excess of 1875 F. that the tensile and/or yield strength of the latter heats is equal to or inferior to the same properties of the steel of Heat 79650 which contains at least delta ferrite. Thus, it is clear that it is only through the use of the process of this invention that it is possible to obtain or develop the full mechanical property potential of the steels substantially free of delta ferrite. This is substantiated more clearly by reference to Heat 9X635 which is even more stable as respects delta ferrite than, for example, Heat 35558. Heat 9X635, when annealed at 1850 F., exhibited a tensile strength of 239,750 p.s.i. whereas after annealing at 1950 F., the tensile strength dropped to 228,780. Similar results were recorded for the yield strength.

All of the foregoing test results were obtained from steels of the transformation hardening type. Reference is now directed to Heats FG-12, FG-ll, FG-li) and Flt-21 which are also of the transformation hardening type, but also have a precipitation hardening mechanism superimposed thereon.

Heat FG-12 contains between 10% and 30% delta ferrite whereas Heats FG-ll, FG-10 and Fl-Zl are substantially free of delta ferrite. From the test results recorded in Table 111, it is seen that 1875 F. is the critical temperature which must not be exceeded in order to derive the beneficial eiiiects of the balance alloy composition to thereby obtain the requisite strength in the alloy. The hardening treatment to which steels FG-12, FG-dl, PG- 10 and FJ-21 were subjected, varied slightly from the hardening treatment set forth hereinbefore. After anhealing at the indicated temperatures, the steels were hardened by trigger annealing at 1750 F. for minutes and thereafter quenched. The steels were then sub-Zero cooled to l00 F. for 8 hours and thereafter tempered at 950 F. for 1 hour. Heat FG-12 contains about 10% delta ferrite and thus is similar to Heat 79650 which contains about 10% delta ferrite. From the test results recorded in Table III for Heat FG-12 and which has 10% delta ferrite, it is apparent that the annealing temperature has little effect on the mechanical properties, especially the tensile strength and the yield strength. Increasing the annealing temperature. from 1800 F. to 1950 F. is efiec- 10 tive for decreasing the tensile strength from 225,000 p.s.i. to 222,000 p.s.i. A similar result was also obtained for the yield strength. Heats FG-l 1; and FG-10 when annealed at the temperature indicated in Table III show a severe drop in both the tensile strength and the yield strength when the temperature of about 1875 F. is exceeded. Thus, in FG-ll, the tensile strength has dropped from 241,000 p.s.i. to 231,000 p.s.i. and the yield strength has dropped from 212,000 p.s.i. to 199,000 p.s.i. when these materials were annealed at 1800 F. and at 1950 F., respectively. The drop is even sharper when the test results recorded in Table III are compared for Heat FG10. The test results for Heat Flt-21 more clearly demonstrate the effect of annealing at temperatures in excess of the upper limit of the critical annealing temperature range.

From the foregoing, it is apparent that the treatment of this invention is effective when used on steels of the transformation hardening type, either with or without a precipitation hardening mechanism superimposed thereon. It is also significant to point out that for the heats listed in Table III, the ductilities have remained substantially constant at about 10% elongation in a 2-inch gauge length. Thus, regardless of the fact that the steel contains delta ferrite or is substantially free of delta ferrite, the process of this invention is effective for producing an outstanding increase in strength of the steels which are substantially free of delta ferrite and without adversely affecting the ductility as measured by the percentage elongation. Note in particular that the steels containing delta ferrite have an elongation of the same order and magnitude as those substantially free of delta ferrite. Thus, it is clear that it is only through treating the steels substantially free of delta ferrite by the process of the present invention that it is possible to obtain an outstanding combination of mechanical properties without adversely afifecting the ductility of the steels.

There are no special skills required to utilize the process of the present invention. All equipment necessary in the commercial utilization of this process is standard steel mill equipment and the temperatures and controls necessary to adequately control the process are normal. The process is outstanding, producing a steel substantially free of delta ferrite having an outstanding balance of mechanical properties contained therein.

This is a continuation-in-part of application Serial No. 840,876, filed September 18, 1959.

We claim:

1. In the process of treating austenitic stainless steels, the steps comprising, selecting an austenitic stainless from the group consisting of austenitic stainless steels of the transformation hardening type and of the transformation and precipitation hardening type and containing from 0.01% to 0.40% carbon, from 0.05% to 8.0% manganese, from 0.05 to 2.0% silicon, from 8.0% to 20.0% chromium, from 1.0% to 13.0% nickel, up to 4.0% molybdenum, up to 0.60% nitrogen, up to 4.0% of metal selected from the group consisting of aluminum, vanadium and copper and the balance iron with the usual impurities and which has a balanced composition of nickel equivalents and chromium equivalents in amounts substantially in accordance with the area aBCdEFGa of the accompanying diagram and which is substantially free of delta ferrite in the annealed condition at room temperature, hot rolling the stainless steel into the form of strip, cleaning the strip, annealing the steel strip at a temperature in excess of 1900 F. to elfect a substantially completesolution of the carbides within the austenite to thereby stabilize the austenite, cooling the steel at a rate suificiently fast to prevent carbide precipitation, cold working the steel in one or moreoperations to within 10% to 40% of the cross sectional area of the finish gauge, reannealing the steel of intermediate gauge at a temperature in excess of 1900 F., cold working the steel to effect a reduction in cross sectional area of atleast 10% to finish gauge to transform a portion of the austenite to martensite and provide the austenite phase with slip bands within the microstructure, continuously critically annealing the steel at a temperature in the range between 1800 F. and 1875 F., cooling the steel to room temperature, trigger annealing the steel at a temperature in the range between 1650 F. and 1750 F. to precipitate a portion of the carbides in a random distribution within the austenite phase, cooling the steel to room temperature at a rate sufliciently fast toprevent further carbide precipitation, sub-zero cooling the steel to a temperature below-80? F. to transformthe austenite to martensite, and thereafter tempering the steel at a temperature in the range between 750 F. and 1050 F.

2. In the process of treating austenitic stainless steels,

the steps comprising, selecting an austenitic stainless from the group consisting of austenitic stainless steels of the transformation hardening type and of the transformation and precipitation-hardening type and containing from 0.01% to 0.40% carbon, from 0.05% to 8.0% manganese, from 0.05% to 2.0% silicon, from 8.0% to 20.0% chromium, from 1.0% to 13.0%, nickel, up to 4.0% molybdenum, up to 0.60% nitrogen, up to 4.0% of metal selected from the group consisting of aluminum, vanadium and copper and the balance iron with the usual impurities and which has a balanced composition of nickel equivalents and chromium equivalents in amounts substantially in accordance with the area aBCdEFGa of the accompanying diagram and which is substantially free of delta ferrite in the annealed condition at room temperature, hot working the stainless steel into the form of strip, cleaning the strip, annealing the steel strip at a temperature in excess of 1900 F. to effect a substantially complete solution of the carbides within the austenite, cooling the steel at a rate suificiently fast to prevent carbide precipitation, cold working the steel to efiect at least a 10% reduction in cross sectional area to finish gauge to transform a portion of the austenite to martensite, critically annealing the steel at a temperature in the range between 1800 F. and 1875 F. for a time period ranging between 15 seconds and 20 minutes depending upon the gauge, trigger annealing the steel at a temperature in the range between 1650 F. and

1750 F. for a time'period of up to about'l hour to precipitate a portion of the carbides in a random distribution throughout the austenite, cooling the steel to roomtemperature at a rate sufliciently fast to prevent further carbide precipitation, sub-zero cooling the steel to a- 1 temperature below 80 F., and thereafter tempering the steel at a temperature in the range between 750 F. and 1050 F. for a time period of at least 1 hour.

3. In the. process of treating austenitic stainless steels,

J the steps'comprising, selectingan austenitic stainless from the group consisting of austenitic stainless steels of the panying diagram and which is substantially free of delta;

ferrite in theannealed condition at room temperature, hot rolling the stainless steel into the form of strip, cleaning the strip, annealing the, steel strip at a temperature in-- excess of 1900 F. to'effecta substantially complete solution of the carbides Within the austenite to thereby stabilize the austenite, cooling the steel at a rate sufficiently fast to prevent carbide precipitation, cold rolling the'steel' in one or more cold rolling operations to an intermediate gauge of within 10% and 40% of the cross sectional area of the finish gauge, reannealing the steel of intermediate gauge at a temperature in excess of 1900 F.,

" puritiesand which has a balanced composition of nickel 18 cooling the steel at a rate sufliciently fast toprevent carbide precipitation, cold rolling the steel-to eifect at least a 10% reduction in cross sectional area to finish gauge to transform a portion of the austenite to martensite and provide the microstructure with slip planes, substantially continuously heating the steel to the critical annealing temperature range of the steel at between 1800 F. and 1875 F. in a time period ranging between 15 seconds and 20 minutes depending upon the gaugeto precipitate a portion of the carbides in a random distribution throughout the microstructure, recrystallize the steel, redissolve a portion of the precipitated carbides and stabilize the austenite phase, cooling the steel at a rate suflicieptly fast to prevent precipitation of the carbides, trigger annealing the steel at a temperature in the range between 1650 F. and 1750 F. for a time period of up to about 1' hour to precipitate a greater portion of the carbides in a random distribution throughout the austenite and to unbalance the austenite phase, cooling the steel to room temperature at a rate sufliciently fast to prevent further carbide precipitation, sub-zero coolingthe steel to' a temperature below F. to transform the austenite to martensite, and thereafter tempering the steel at a temperature in the range between 750 F. and 1050 F.

4. In theprocess of treating austenitic stainless steels, the steps comprising, selectingan austenitic stainless steel from the group consisting of austenitic stainless steels of the transformation hardening type and of the transformation and precipitation hardening type and containing from 0.01% to 0.40% carbon, from 0.05% to 8.0% manganese, from 0.05% to 2.0% silicon, from 8.0% to 20.0% chromium, from 1.0% to 13.0% nickel, up to 4.0% molybdenum, up to 0.60% nitrogen, up to 4.0% of metal selected from the group consisting of aluminum,vanadium and copper and the balance iron with the usual impurities and which has a balanced composition of nickel equivalents and chromium equivalents in amounts substantially in accordance with the area ABCDEFGA of the accompanying diagram and, which contains not more than 5% delta-ferrite in the annealed condition at room temperature, hot working the stainless steel into the form of strip, cleaning the strip, annealing the steel strip at a temperature in excess of 1900 F. to elfect complete solution of the carbides within the austenite, cooling the steel at arate sufiicently fast to prevent carbide precipitation,

" cold working the steel to finish gauge to transform a portion of the austenite 'to martensite, critically annealing the steel at a temperature in the range between 1800" F. and 1875 F. to precipitate a portion of the carbides in a random distribution throughout the microstructure, trigger annealing the steel at a temperature in the range between 1650 F. and 1750 F. to unbalance the austenite and precipitate additional carbides throughout the microthe steps comprising, selecting an austenitic stainless steel from the group consisting of austenitic stainless steels of the transformation hardening type and of the transformation and precipitation hardening typeand containing from 0.01% to 0.40% carbon, from 0.05 to 8.0% manganese, from 0.05 to 2.0% silicon, from 8.0% to 20.0% chromium, 'from 1.0% to 13.0% nickel, up to 4.0% molybdenum, up-to 0.60% nitrogen, up to 4.0% of metal selected from the group consisting of aluminum, vanadiumand copper andthe balance iron with the usual i'rn' equivalents and chromium equivalents in amounts 'substantiallyin accordance with the area aBCdEFGa of'the accompanyingdiagram and whichv is substantially austenitic inthe annealed condition at room temperature, hot

working the stainless steel into the form of strip, cleaning the strip, annealing the steel strip at a temperature in excess of 1900 F. to effect complete solution of the carbides within the austenite, cooling the steel at a rate sufficiently fast to prevent carbide precipitation, cold working the steel to finish gauge totransform a portion of the austenite to martensite, critically annealing the steel at a temperature infthe range between l800'F. and 1875 F. to precipitate a portion of the carbides in a random distribution throughout the microstructure, trigger annealing the steel at a temperature in the range between 1650 F. and 1750 F. to unbalance the austenite and precipitate additional carbides throughout the microstructure, cooling the steel to room temperature at a rate sufliciently fast to prevent further carbide precipitation, sub-zero cooling the steel to a temperature below 80 F. to transform the austenite to martensite, and thereafter tempering the steel at a temperature in the range between 750 F. and 1050 F.

6. In the process of treating austenitic stainless steels, the steps comprising, selecting an austenitic stainless from the group consisting of austenitic stainless steels of the transformation hardening type and of the transformation and precipitation hardening type and containing from 0.01% to 0.40% carbon, from 0.05 to 8.0% manganese, from 0.05% to 2.0% silicon, from 8.0% to 20.0% chromium, from 1.0% to 13.0% nickel, up to 4.0% molybdenum, up to 0.60% nitrogen, up to 4.0% of metal selected from the group consisting of aluminum, vanadium and copper and the balance iron with the usual impurities and which has a balanced composition of nickel equivalents and chromium equivalents in amounts substantially in accordance with the area ABCDEFGA of the accompanying diagram and which contains a maximum of delta ferrite in the annealed condition at room temperature, hot rolling the stainless steel into the form of strip, cleaning the strip, annealing the steel strip at a temperature in-excess of 1900" F. to effect a substantially complete solution of the carbides within the austenite to thereby stabilize the austenite, cooling the steel at a rate sutficiently fast to prevent carbide precipitation, cold working the steel in one or more operations to within 10% and 40% of the cross sectional area of the finish gauge, reannealing the steel of intermediate gauge at a temperature in excess of 1900 F., cold working the steelv portion of the carbides in a random distribution within the austenite phase, cooling the steel to room temperature at a rate sufficiently fast to prevent further carbide precipitation, sub-zero cooling the steel to a temperature below 80 F. to transform the austenite to martensite, and thereafter tempering the steel at a temperature in the range between750 F. and 1050'F.

7. In the process of treating austenitic stainless steels, the step comprising, selecting an austenitic stainless from r 20 s A diagram and which contains a maximum of 5%delta ferrite in the annealed condition at roomtemperature, hot working the stainless steel into the form of strip, cleaning the strip, annealing the steel strip at a temperature in excess of l900 F.'to effect a substantially complete solution of the carbides within the austenite, cooling the steel at a ratesufliciently fast to prevent carbide precipitation, cold working the steel to effect at least a 10% reduction in cross sectional area to finish gauge to transform a portion of the austenite to martensite, critically annealing the steel at atemperature in the range between 1800? F. and 1875 F. for a time period ranging between 15 seconds and 20 minutes depending upon the gauge,.trigger annealing the steel at a temperature in the range between 1650 F. and 1750 F. for a time period of up to about 1 hour .to precipitate a portion of the carbides in a random distribution throughout the austenite, cooling the steel to room temperature at a rate sufiiciently fast to prevent further carbide precipitation, sub-zero cooling the steel to a temperature below 80 F., and thereafter tempering the steel at a temperature in the range between 750 F. and 1050 F. for a time period of at least 1 hour.

8. In the process of treating austenitic stainless steels, the steps comprising, selecting an austenitic stainless from the group consisting of. austenitic stainless steels of the transformation hardening type and of the transformation and precipitation hardening type and containing from 0.01% to 0.40% carbon, from 0.05% to 8.0% manganese,

from 0.05% to 2.0% silicon, from 8.0% to 20.0% chromium, from 1.0% to 13.0%. nickel, up to 4.0% molybdenum, up to 0.60% nitrogen, up to 4.0% of metal selectedfrom the group consisting ofaluminum, vanadium and copper and the balance iron with the usual impurities and which has a balanced composition of nickel equivalents and chromium equivalents in amounts substantially in accordance with the area ABCDEFGA of the accompanying diagram and which contains a maximum of 5% delta ferrite in the annealed condition at room temperature, hot rolling the stainless steel into the form of strip, cleaning the strip, annealing the steel strip at a temperature in excess of 1900 F. to effect a substantially complete solution of the carbides within the austenite to thereby stabilize the austenite, cooling the steel at a rate 'sufiiciently fast to prevent carbide precipitation, cold roll- .ing the steel in one or more cold rolling operations to an intermediate gauge of within 10% and 40% of the cross sectional area of the finish gauge, reannealing the steel of intermediate gauge at a temperaturein excess of 1900 1 F-., cooling the steel at a rate sufficiently fast to prevent the group consisting of austenitic stainless steels of the transformation hardening type and of the transformation and precipitation hardening type and containing from 0.01% to0.40% carbon, from 0.05 to 8.0% manganese, from 0.05 to 2.0% silicon, from 8.0% to 20.0% chro mium, from 1%.to 13% nickel, up to 4.0% molybdenum, up'to 0.60% nitrogen, up to 4.0% of metal'selected from the group consisting of aluminum, vanadium and copper and the balance iron with the usual impurities and which has a balanced composition of nickel equivalents and chromium equivalents in amounts substantially in accordance with the area ABCDEFGA of the'accompanying carbide precipitation, cold rolling the steelto effect at least a 10% reduction in cross sectional area to finish gauge to transform a portion of the austenite ..to martensite and provide the microstructure with slip planes, substantially continuously heating the steel'to the critical anealing temof the precipitated carbides and stabilize the austenite Tphase, cooling the steel at a rate sufficiently fast to prevent precipitation of the carbides, trigger annealingthe steel at --a temperature in the range between 1650 F. and 1750 F. for a time period of up to about 1 hour to precipitate a greater portion of the earbidesin a random distribution sub-zero cooling the steel to a temperature below F. to transform the austenite to martensite, and thereafter tempering the steel at a temperature in the range between 7509a and 1050 (References on following page) 21 22 References Cited in the file of this patent OTHER REFERENCES 7 UNITED STATES PATENTS ARMCO, Product Data Bulletin, February 1, 1959, 2,799,602 Lena -i July 16, 1957 Armco PH 15-7 Mo Sheet, Strip and Plate, Market De- 2320708 Waxweiler 1953 velopment Division, Armco Steel Corporation, Middle- 2,894,867 Smith July 14, 1959 l 2,958,618 Allen Nov. 1, 1960 town, Ohm (p- 243 rehed 9p 

1. IN TH PROCESS OF TREATING AUSTENITIC STAINLESS STEELS, THE STEPS COMPRISING, SELECTING AN AUSTENITIC STAINLESS FROM THE GROUP CONSISTING OF AUSTENITIC STAINLESS STEELS OF THE TRANSFORMATION HARDENING TYPE AND OF THE TRANSFORMATION AND PRECIPITATION HARDENING TYPE AND CONTAINING FROM 0.1% TO 0.40% CARBON, FROM 0.05% TO 8.0% TO 20.0% CHROMIUM, FROM 1.0% TO 13.0% NICKEL, UP TO 4.0% MOLYBDENUM, UP TO 0.60% NITROGEN, UP TO 4.0% METAL SELECTED FROM THE GROUP CONSISTING OF ALUMINUM, VANADIUM AND COPPER AND THE BALANCE IRON WITH THE USUAL IMPURITIES AND WHICH HAS A BALANCED COMPOSITION OF NICKEL EQUIVALENTS AND CHROMIUM EQUIVALENTS IN AMOUNTS SUBSTANTIALLY IN ACCORDANCE WITH THE AREA OF ABCDEFGA OF THE ACCOMPANYING DIAGRAM AND WHICH IS SUBSTANTIALLY FREE OF DELTA FERRITE IN THE ANNEALED CONDITION AT ROOM TEMPERATURE, HOT ROLLING THE STAINLESS STEEL INTO THE FORM OF STRIP, CLEANING THE STRIP, ANNEALING THE STEEL STRIP AT A TEMPERATURE IN EXCESS OF 1900*F. TO EFFECT A SUBSTANTIALLY COMPLETE SOLUTION OF THE CARBIDES WITHIN THE AUSTENITE TO THEREBY STABILIZE THE AUSTENITE, COOLING THE STEEL AT A RATE SUFFICIENTLY FAST TO PREVENT CARBIDE PRECIPITATION, COLD WORKING THE STEEL IN ONE OR MORE OPERATIONS TO WITHIN 10% TO 40% OF THE CROSS SECTIONAL AREA OF THE FINISH GAUGE, REANNEALING THE STEEL OF INTERMEDIATE GAUGE AT A TEMPERATURE IN EXCESS OF 1900*F., COLD WORKING THE STEEL TO EFFECT A REDUCTING IN CROSS SECTIONAL AREA OF AT LEAST 10% TO FINISH GAUGE TO TRANSFORM A PORTION OF THE AUSTENITE TO MARTENSITE AND PROVIDE THE AUSTENITE PHASE WITH SLIP BANDS WITHIN THE MICROSTRUCTURE, CONTINUOUSLY CRITICALLY ANNEALING THE STEEL AT A TEMPERATURE IN THE RANGE BETWEEN 1800*F. AND 1875*F., COOLING THE STEEL TO ROOM TEMPERATURE, TRIGGER ANNEALING THE STEEL AT A TEMPERATURE IN THE RANGE BETWEEN 1650*F. AND 1750* F. TO PRECIPITATE A PORTION OF THE CARBIDES IN A RANDOM DISTRIBUTION WITHIN THE AUSTENITE PHASE, COOLING THE STEEL TO ROOM TEMPERATURE AT A RATE SUFFICIENTLY FAST TO PREVENT FURTHER CARBIDE PRECIPITATION, SUB-ZERO COLLING THE STEEL TO A TEMPERATURE BELOW -80*F. TO TRANSFORM THE AUSTENITE TO MARTENSITE, AND THEREAFTER TEMPERING THE STELL AT A TEMPERATURE IN THE RANGE BETWEEN 750*F. AND 1050*F. 